KR101598307B1 - High-strength steel sheet having superior impact resistance, method for producing same, high-strength galvanized steel sheet, and method for producing same - Google Patents

Abstract

The steel sheet according to any one of claims 1 to 3, wherein the steel sheet contains C, Si, Mn, P, S, Al, Ti, N and O at a predetermined content and the balance contains iron and unavoidable impurities. / 8 thickness to 3/8 thickness, wherein the retained austenite has an average aspect ratio of not more than 2.0, and the amount of Mn in the retained austenite is an average Mn of at least 1.1 times and having an average particle diameter of 0.5 占 퐉 or less and has a density of 1.0N / mm2 or less of AlN grains having an average particle diameter of 1 占 퐉 or more and has an excellent impact resistance characteristic with a tensile maximum strength of 900 MPa or more High strength steel plate.

Description

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a high strength steel sheet excellent in impact resistance, a method of manufacturing the same, a high strength galvanized steel sheet and a method of manufacturing the same. BACKGROUND ART [0002]

The present invention relates to a high strength steel sheet, a method of manufacturing the same, a high strength galvanized steel sheet and a manufacturing method thereof, and more particularly to a high strength steel sheet having excellent impact resistance characteristics and a manufacturing method thereof. This application is based upon and claims the benefit of priority from Japanese Patent Application No. 2011-167661 filed on July 29, 2011, the contents of which are incorporated herein by reference.

In recent years, in order to increase the collision safety while reducing the weight of the automobile, it is required to improve the strength of the steel sheet used for automobiles and to improve the impact resistance characteristics.

Patent Document 1 discloses a high strength steel sheet having a large impact absorbing energy. The steel sheet comprises 0.05 to 0.3% of C, 2.0% or less of Si, 0.01 to 2.0% of Al, 0.5 to 4.0% of Mn, %, P: not more than 0.1%, S: not more than 0.1%, N: not more than 0.01%, the balance being Fe and inevitable impurities, and 1.5-3.0 × C≤Si + Al? 3.5-5.0 × C and Mn + Ni / 3) ≥ 1.0 (%), and the baking hardening amount of the steel sheet is 50 MPa or more.

Patent Document 2 discloses a high tensile strength steel sheet excellent in collision absorbency and having a volume ratio VB given by a formula VB? (TSs / 60) -1 (TSs: tensile strength (MPa) in a static tensile test) , A retained austenite having a C content of 1.2% by mass or less and a volume ratio of 5% or more, and a residual amount of ferrite, wherein the yield ratio in the static tensile test is 0.6 or more, TSs / TSs > 0.8 + (300 / TSs) (TSd: tensile strength in a dynamic tensile test at a deformation rate of 1000 / s) of the tensile strength in the tensile test (MPa)) of a high-ductility high-tensile steel sheet.

Patent Document 3 discloses a method for producing a high-strength cold-rolled steel sheet having excellent impact properties, which comprises 0.08 to 0.18 mass% of C, 1.00 to 2.0 mass% of Si, 1.5 to 3.0 mass% of Mn, , A S content of 0.005 mass% or less and a T.Al content of 0.01 to 0.1 mass% and a Mn segregation degree defined by a formula (Mn segregation degree = (Mn concentration in slab center - base Mn concentration) / base Mn concentration) 1.10 slabs are hot rolled and then cold rolled and then heated in a continuous annealing line in a two phase or single phase region at 750 to 870 占 폚 for a holding time of 60 seconds or more and then cooled to a mean temperature of 720 to 600 占 폚 After cooling to a speed of 10 ° C / s or less, cooling to 350 to 460 ° C at an average cooling rate of 10 ° C / s or more and holding for 30 seconds to 20 minutes, cooling to room temperature was performed to obtain a polycrystalline ferrite + acicular ferrite + bainite + Residual austenite + martensite A five-phase structure is described.

Patent Document 4 discloses a steel sheet for use in a steel sheet for automobiles, which comprises 0.05 to 0.25% of C, 0.5% or less of Si, 1 to 3% of Mn, 0.1% or less of P, Al: 0.1 to 2%, N: and containing less than 0.005%, and Si + Al≥0.6%, (0.0006Al) % ≤N≤0.0058% - (0.0026 × Al)%, Al≤ (1.25 × C 0 .5 - 0.57Si + 0.625Mn)%, and the remaining amount of Fe and unavoidable impurities are contained in the hot-dip galvanized steel sheet.

Patent Document 5 discloses a high strength alloyed hot-dip galvanized steel sheet excellent in energy absorption characteristics, which contains 0.05 to 0.20 mass% of C, 0.3 to 1.5 mass% of Si, 1.0 to 2.5 mass% of Mn and 0.1 mass% And a balance of Fe and unavoidable impurities, and a total of 25 to 50% by volume of one or two of martensite and retained austenite, and the remainder being composed of ferrite and bainite It is described that a steel sheet having a microstructure is used as a substrate and both surfaces of the steel sheet are subjected to alloying hot dip galvanizing.

Patent Document 6 discloses a high-strength, high-strength cold-rolled steel sheet excellent in surface properties and impact absorbability, which comprises 0.06 to 0.25% of C, 2.5 to 3.0% of Si, 0.5 to 3.0% of Mn, , Ti: 0.003 to 0.08%, N: 0.01% or less, and the balance of Fe and inevitable impurities, wherein the Ti content is (48/14) N (48/14) N + (48/32) S + 0.01, and the structure after the cold-rolling-recrystallization annealing is a structure containing residual austenite having a volume ratio of 5% or more.

Patent Document 7 discloses a high ductile high strength steel sheet excellent in low temperature toughness, having a structure in which the area% is 60% or more of bainite, 1 to 20% of residual? And the remaining portion is substantially ferrite, Are present in the bainite particles.

[Patent Document 1]: Japanese Patent Application Laid-Open No. 2001-11565 [Patent Document 2]: Japanese Patent Application Laid-Open No. 2002-294400 [Patent Document 3]: Japanese Patent Application Laid-Open No. 2004-300452 [Patent Document 4]: Japanese Patent Application Laid-Open No. 2006-307327 [Patent Document 5]: Japanese Patent Application Laid-Open No. 2009-68039 [Patent Document 6]: JP-A-10-130776 [Patent Document 7]: JP-A-11-21653

However, in the conventional technology, sufficient impact resistance characteristics can not be obtained in a steel sheet having a high tensile strength of 900 MPa or more, and furthermore, it is required to further improve the impact resistance characteristics.

In view of the above circumstances, the present invention provides a high strength steel sheet having excellent impact resistance characteristics, a method for producing the same, and a high strength galvanized steel sheet formed by forming a zinc plating layer on the surface of a high strength steel sheet excellent in impact resistance .

The inventors of the present invention have conducted intensive investigations to obtain a high strength steel sheet having a maximum tensile strength of 900 MPa or more in which excellent impact resistance characteristics can be obtained. As a result, the present inventors have found that the steel sheet has a predetermined chemical composition containing 0.001 to 0.050% of Al, 0.0010 to 0.0150% of Ti, and 0.0001 to 0.0050% of N, Wherein the steel sheet contains 1 to 8% of retained austenite in a volume fraction of 1/8 to 3/8 of the thickness, the average aspect ratio of the retained austenite is 2.0 or less, And the density of the AlN grains having a grain size of 1 탆 or more is 1.0 / mm 2 or less, which contains TiN grains having an average grain size of 0.5 탆 or less.

That is, the high-strength steel sheet contains Al, Ti and N within the above-mentioned range, and the generation of fine TiN grains having an average grain diameter of 0.5 탆 or less produces AlN grains having an average grain diameter of 1 탆 or more The density of the AlN particles having a particle diameter of 1 mu m or more is less than 1.0 / mm < 2 >. As a result, the high-strength steel sheet is prevented from destruction, which is the starting point of the AlN particles.

In the high-strength steel sheet, the volume fraction of retained austenite as the fracture origin is as low as 1 to 8%, the retained austenite has a stable shape with excellent isotropy with an average aspect ratio of 2.0 or less, Is at least 1.1 times the average Mn amount. Therefore, in the high-strength steel sheet, breakage of the retained austenite as a starting point is prevented.

As described above, in the high-strength steel sheet, breakage that is the starting point of the AIN particles and breakage that is the starting point of the retained austenite are prevented, and therefore excellent impact resistance characteristics can be obtained.

The present invention has been completed based on this knowledge, and the essential points thereof are as follows.

(One)

The steel sheet according to any one of claims 1 to 3, wherein the steel sheet contains 0.075 to 0.300% of C, 0.30 to 2.50% of Si, 1.30 to 3.50% of Mn, 0.001 to 0.050% of P, 0.0001 to 0.0050% of S, 0.001 to 0.050% of Al, And a ratio of O to 0.0001 to 0.0030%, the balance being iron and inevitable impurities, and having a thickness of 1/8 to 3/4, 8 to 8% by volume of the retained austenite, wherein the retained austenite has an average aspect ratio of not more than 2.0, and the amount of dissolved Mn in the retained austenite is not less than 1.1 times , And has a steel sheet structure containing TiN particles having an average particle diameter of 0.5 占 퐉 or less and a density of AlN particles having a particle diameter of 1 占 퐉 or more of 1.0 pieces / mm2 or less and having a maximum tensile strength of 900 MPa or more.

(2)

Wherein the steel sheet structure contains ferrite having a volume fraction of 10 to 75% or less and either or both of bainitic ferrite and bainite in a total of 10 to 50% and 10 to 50% or less of tempering martensite,

The high strength steel sheet according to (1), wherein the pearlite has a volume fraction of 5% or less and fresh martensite has a volume fraction of 15% or less.

(3)

The high strength steel sheet according to (1), further comprising at least one of Nb: 0.0010 to 0.0150%, V: 0.010 to 0.150%, and B: 0.0001 to 0.0100% in mass%.

(4)

The steel sheet according to any one of the items (1) to (3), further comprising one or two or more of Cr: 0.01 to 2.00%, Ni: 0.01 to 2.00%, Cu: 0.01 to 2.00%, Mo: 0.01 to 1.00% , And (1).

(5)

The high strength steel sheet according to (1), further comprising one or more kinds of Ca, Ce, Mg, Zr, Hf and REM in a total amount of 0.0001 to 0.5000 mass%.

(6)

A high strength galvanized steel sheet excellent in impact resistance as described in (1), wherein a zinc plated layer is formed on the surface.

(7)

A high strength galvanized steel sheet excellent in impact resistance as described in (6), wherein a coating comprising a composite oxide containing phosphorous oxide and / or phosphorus is formed on the surface of the zinc plated layer.

(8)

The steel sheet according to any one of claims 1 to 3, wherein the steel sheet contains 0.075 to 0.300% of C, 0.30 to 2.50% of Si, 1.30 to 3.50% of Mn, 0.001 to 0.050% of P, 0.0001 to 0.0050% of S, 0.001 to 0.050% of Al, A slab comprising 0.0150 to 0.0150% of N, 0.0001 to 0.0050% of O, 0.0001 to 0.0030% of O and the balance of iron and inevitable impurities is heated to 1210 캜 or higher, and in a temperature range of at least 1100 to 1000 캜 , The rolling is carried out under a condition satisfying the following expression (1), and the rolling is completed at a hot rolling temperature of not less than a temperature higher than 800 ° C and an Ar ₃ transformation point of not more than 970 ° C, A cold rolling step of cold rolling at a reduction ratio of 30 to 75% after the hot rolling step and a cold rolling step of cooling the hot rolling at a temperature of 550 to 700 DEG C Heating the range to an average heating rate of 10 [deg.] C / sec or less Cooling at an average cooling rate of 1.0 to 10.0 ° C / sec in the temperature range from the maximum heating temperature to 700 ° C, and the temperature of 700 ° C to 500 ° C in the range of the maximum heating temperature (Ac ₁ transformation point +40) And a continuous annealing step of performing annealing by cooling at an average cooling rate of 5.0 to 200.0 캜 / second and a rectifying treatment at a temperature range of 350 to 450 캜 for 30 to 1000 seconds.

Ti represents the machining temperature of the i-th pass, ti represents the elapsed time from the i-th pass to the (i + 1) th pass, and ei represents the reduction ratio of the i-th pass.

(9)

A galvanized steel sheet produced by the method of manufacturing a galvanized steel sheet having excellent impact resistance in a continuous annealing process according to the production process described in (8), wherein zinc plating is applied to the surface of the steel sheet after the rectifying treatment to form a galvanized layer.

(10)

(8), before or after the rectifying treatment in the temperature range of 350 to 450 占 폚 after cooling in the temperature range of 700 to 500 占 폚, the steel sheet is subjected to the galvanizing bath To form a galvanized layer on the surface of the steel sheet, wherein the galvanized steel sheet is excellent in impact resistance.

(11)

A process for producing a high strength galvanized steel sheet excellent in impact resistance as claimed in (10), wherein the steel sheet is reheated to 460 to 600 ° C after immersed in the zinc plating bath and maintained for 2 seconds or more to alloy the galvanized layer.

(12)

(10), wherein a coating film comprising a composite oxide containing phosphorus oxide and / or phosphorus is provided on the surface of the zinc plated layer after the zinc plating layer is formed, A method of manufacturing a steel sheet.

(13)

(11), wherein the zinc plating layer is alloyed and a coating film comprising a composite oxide containing phosphorus oxide and / or phosphorus is provided on the surface of the alloyed zinc plating layer, A method for producing a galvanized steel sheet.

In the high strength steel sheet of the present invention, since the AlN grains and retained austenite are prevented from becoming destructive starting points, a high strength steel sheet having excellent impact resistance characteristics and having a tensile maximum strength of 900 MPa or more can be obtained. Further, according to the method for producing a high strength steel sheet of the present invention, it is possible to provide a high strength steel sheet having an excellent impact resistance characteristic and a tensile maximum strength of 900 MPa or more. Further, according to the present invention, there can be provided a high strength galvanized steel sheet in which a galvanized layer is formed on the surface of a high-strength steel sheet excellent in impact resistance characteristics and a method for producing the same.

(Chemical composition)

First, the chemical composition (composition) of the high-strength steel sheet of the present invention will be described. In the following description, [%] is [mass%].

The high strength steel sheet of the present invention is characterized in that the steel sheet contains 0.075 to 0.300% of C, 0.30 to 2.50% of Si, 1.30 to 3.50% of Mn, 0.001 to 0.050% of P, 0.0001 to 0.0050% of S, 0.001 to 0.050% : 0.0010 to 0.0150%, N: 0.0001 to 0.0050%, and O: 0.0001 to 0.0030%, with the balance being iron and inevitable impurities.

&Quot; C: 0.075 to 0.300% "

C is contained in order to increase the strength of the high strength steel sheet. However, if the content of C exceeds 0.300%, the weldability becomes insufficient. From the viewpoint of weldability, the content of C is preferably 0.250% or less, and more preferably 0.220% or less. On the other hand, if the content of C is less than 0.075%, the strength is lowered and the maximum tensile strength of 900 MPa or more can not be secured. In order to increase the strength, the content of C is preferably 0.090% or more, more preferably 0.100% or more.

&Quot; Si: 0.30 to 2.50% "

Si is an element required for suppressing the generation of iron-based carbides in the steel sheet and for increasing the strength and moldability. However, if the content of Si exceeds 2.50%, the steel sheet becomes brittle and ductility deteriorates. From the viewpoint of ductility, the content of Si is preferably 2.20% or less, more preferably 2.00% or less. On the other hand, when the content of Si is less than 0.30%, a large amount of coarse iron carbide is generated in the annealing step, and the strength and formability are deteriorated. From this viewpoint, the lower limit value of Si is preferably 0.50% or more, more preferably 0.70% or more.

"Mn: 1.30 to 3.50%"

Mn is added to the steel sheet of the present invention to increase the strength of the steel sheet. However, if the content of Mn exceeds 3.50%, a coarse Mn-enriched portion occurs at the central portion of the plate thickness of the steel sheet, brittleness tends to occur, and troubles such as breaking of the casted slab are liable to occur. If the content of Mn exceeds 3.50%, the weldability also deteriorates. Therefore, the content of Mn should be 3.50% or less. From the viewpoint of weldability, the content of Mn is preferably 3.20% or less, more preferably 3.00% or less. On the other hand, if the content of Mn is less than 1.30%, a large amount of soft structure is formed during cooling after annealing, and it becomes difficult to secure a tensile maximum strength of 900 MPa or more. From this point of view, it is necessary to set the Mn content to 1.30% or more. In order to increase the strength, the content of Mn is preferably 1.50% or more, more preferably 1.70% or more.

"P: 0.001 to 0.050%"

P tends to segregate at the central portion of the plate thickness of the steel sheet, and brittle the welded portion. If the content of P exceeds 0.050%, the welded portion is significantly brittle, so that the content of P is limited to 0.050% or less. The effect of the present invention is exhibited even if the lower limit of the P content is not specifically defined. However, setting the P content to less than 0.001% involves a significant increase in the production cost, so 0.001% is the lower limit.

"S: 0.0001 to 0.0050%"

S adversely affects the weldability and the manufacturability during casting and hot rolling. In addition, since it forms a sulfide in association with Ti to inhibit Ti from becoming a nitride and indirectly cause the formation of Al nitride, the upper limit of the S content is set to 0.0050%. From this viewpoint, the content of S is preferably 0.035% or less, and more preferably 0.0025% or less. The effect of the present invention is exhibited even if the lower limit of the S content is not specially defined. However, if the S content is less than 0.0001%, the manufacturing cost is significantly increased, so the lower limit is 0.0001%.

&Quot; Al: 0.001% to 0.050% "

When Al is added in a large amount, a coarse nitride is formed to lower the elongation at low temperature and lower the impact resistance, so that the upper limit of Al content is set to 0.050%. In order to avoid generation of a coarse nitride, the content of Al is preferably 0.035% or less. The effect of the present invention is exhibited even if the lower limit of the Al content is not specially defined. However, if the Al content is less than 0.001%, the production cost is significantly increased, so that the lower limit is 0.001%. Further, Al is an effective element as a deacidification material, and from this viewpoint, the content of Al is preferably 0.005% or more, more preferably 0.010% or more.

"N: 0.0001 to 0.0050%"

N forms a coarse nitride which is a starting point of fracture at a low temperature to deteriorate the impact resistance, so that it is necessary to suppress the addition amount. When the content of N exceeds 0.0050%, the influence becomes remarkable, so the range of the N content is 0.0050% or less. From this viewpoint, the content of N is preferably 0.0040% or less, and more preferably 0.0030% or less. When the content of N is less than 0.0001%, the production cost is greatly increased, so that the lower limit of 0.0001% is set as the lower limit.

"O: 0.0001 to 0.0030%"

O forms a coarse oxide and generates a fracture origin at a low temperature, so that it is necessary to suppress the content. When the content of O exceeds 0.0030%, the effect becomes significant, so the upper limit of the O content is 0.0030% or less. From this viewpoint, the content of O is preferably 0.0020% or less, more preferably 0.0010% or less. The effect of the present invention is exhibited even if the lower limit of the O content is not specifically defined. However, when the content of O is less than 0.0001%, the production cost is significantly increased, so the lower limit is 0.0001%.

&Quot; Ti: 0.0010 to 0.0150% "

Ti is an element that inhibits the formation of coarse Al nitride by forming a fine nitride by performing hot rolling under suitable conditions, and it reduces the fracture origin at low temperature and improves the impact resistance characteristic. In order to obtain this effect, the Ti content needs to be 0.0010% or more, the Ti content is preferably 0.0030% or more, and more preferably 0.0050% or more. On the other hand, if the content of Ti exceeds 0.0150%, the formation of soft carbonitride in the steel sheet is deteriorated due to the precipitation of fine carbonitride, thereby lowering the elongation at low temperature. Therefore, the content of Ti is set to 0.0150% or less. From the viewpoint of moldability, the content of Ti is preferably 0.0120% or less, and more preferably 0.0100% or less.

The high-strength steel sheet of the present invention may contain the following elements as required.

&Quot; Nb: 0.0010 to 0.0150% "

Nb is an element which inhibits the formation of coarse Al nitride by forming a fine nitride by performing hot rolling under suitable conditions and reduces the fracture origin at low temperatures. To obtain this effect, the content of Nb is preferably 0.0010% or more, more preferably the content of Nb is 0.0030% or more, and more preferably 0.0050% or more. On the other hand, if the content of Nb exceeds 0.0150%, the formability of a soft portion in the steel sheet deteriorates due to precipitation of fine carbonitride, and the reduction of the elongation at low temperature lowers, so that the content of Nb is preferably 0.0150% Do. From the viewpoint of moldability, the content of Nb is more preferably 0.0120% or less, and still more preferably 0.0100% or less.

&Quot; V: 0.010 to 0.150% "

V is an element that inhibits the formation of coarse Al nitride by forming a fine nitride by performing hot rolling under suitable conditions and reduces the fracture origin at low temperature. In order to obtain this effect, the V content needs to be 0.010% or more, and the content is preferably 0.030% or more, and more preferably 0.050% or more. On the other hand, if the content of V exceeds 0.150%, the formability of a soft portion in the steel sheet deteriorates due to the precipitation of fine carbonitride, and the reduction of the shrinkage value at low temperatures lowers, so that the content of V is preferably 0.150% Do. From the viewpoint of moldability, the content of V is more preferably 0.120% or less, and further preferably 0.100% or less.

"B: 0.0001 to 0.0100%"

B is an element which inhibits the formation of coarse Al nitride by forming a fine nitride by performing hot rolling under suitable conditions and reduces the fracture origin at low temperature. In order to obtain this effect, the content of B is preferably 0.0001% or more, and the content of B is preferably 0.0003% or more, and more preferably 0.0005% or more. Further, B is an element effective for increasing the strength by suppressing the phase transformation at a high temperature and may be further added. However, if the B content exceeds 0.0100%, the workability in hot is degraded and the productivity is lowered. % Or less. From the viewpoint of productivity, the content of B is more preferably 0.0050% or less, and still more preferably 0.0030% or less.

&Quot; Cr: 0.01 to 2.00% "

Cr is an element effective for increasing the strength by suppressing phase transformation at a high temperature and may be added instead of a part of C and / or Mn. If the content of Cr exceeds 2.00%, the workability in hot is impaired and the productivity is lowered. Therefore, the content of Cr is preferably 2.00% or less. The effect of the present invention is exhibited even if the lower limit of the Cr content is not specifically defined. However, in order to sufficiently obtain the effect of increasing the strength by Cr, the Cr content is preferably 0.01% or more.

&Quot; Ni: 0.01 to 2.00% "

Ni is an element effective for increasing the strength by suppressing phase transformation at a high temperature, and may be added instead of a part of C and / or Mn. If the content of Ni exceeds 2.00%, the weldability is impaired. Therefore, the Ni content is preferably 2.00% or less. The effect of the present invention is exhibited even if the lower limit of the Ni content is not specifically defined. However, in order to sufficiently obtain the effect of increasing the strength by Ni, the content of Ni is preferably 0.01% or more.

&Quot; Cu: 0.01 to 2.00% "

Cu is a fine particle which increases strength by being present in steel, and can be added instead of C and / or a part of Mn. If the content of Cu exceeds 2.00%, the weldability is impaired. Therefore, the content of Cu is preferably 2.00% or less. The lower limit of the Cu content is not particularly limited, but the effect of the present invention is exhibited. However, in order to sufficiently obtain the effect of increasing the strength by Cu, the content of Cu is preferably 0.01% or more.

&Quot; Mo: 0.01 to 1.00% "

Mo is an element effective for increasing the strength by suppressing phase transformation at a high temperature and may be added instead of C and / or a part of Mn. If the Mo content exceeds 1.00%, the workability in hot work is impaired and the productivity is lowered. From this point of view, the content of Mo is preferably 1.00% or less. The lower limit of the Mo content is not particularly limited, but the effect of the present invention is exhibited. However, in order to sufficiently obtain the effect of increasing the strength by Mo, the Mo content is preferably 0.01% or more.

&Quot; W: 0.01 to 1.00% "

W is an element effective for increasing the strength by suppressing phase transformation at a high temperature, and may be added instead of a part of C and / or Mn. If the content of W exceeds 1.00%, the workability in hot is impaired and the productivity is lowered. Therefore, the content of W is preferably 1.00% or less. The effect of the present invention is exhibited even if the lower limit of the W content is not specifically defined. However, in order to sufficiently obtain the effect of increasing the strength by W, the content of W is preferably 0.01% or more.

One or more of Ca, Ce, Mg, Zr, Hf and REM in a total amount of 0.0001 to 0.5000%

Ca, Ce, Mg, Zr, Hf and REM are effective elements for improving the formability, and one or two or more kinds can be added. However, if the total content of one or more of Ca, Ce, Mg, Zr, Hf and REM exceeds 0.5000%, the ductility may be deteriorated. Therefore, the total content of the respective elements is preferably 0.5000% or less. The effect of the present invention is exhibited even if the lower limit of the content of one or more of Ca, Ce, Mg, Zr, Hf and REM is not particularly specified. However, in order to sufficiently obtain the effect of improving the formability of the steel sheet, Is preferably 0.0001% or more. From the viewpoint of moldability, the total content of one or more of Ca, Ce, Mg, Zr, Hf and REM is more preferably 0.0005% or more, and still more preferably 0.0010% or more.

REM is an abbreviation of Rare Earth Metal and refers to an element belonging to the lanthanide series. In the present invention, REM or Ce is often added by mischmetal, and in addition to La or Ce, a lanthanoid-based element may be contained in combination. The effect of the present invention is exhibited even when the lanthanoid-based element other than La or Ce is contained as an inevitable impurity. Even when the metal La or Ce is added, the effect of the present invention is exhibited.

(Steel plate organization)

The reason for defining the structure of the high strength steel sheet of the present invention is as follows.

&Quot; TiN particles "

The steel sheet structure of the high strength steel sheet of the present invention contains TiN grains having an average grain size of 0.5 탆 or less. The coarse TiN particles become the fracture origin, but the fine TiN particles having an average particle size of 0.5 탆 or less do not act as the fracture origin. The average particle diameter of the TiN particles is preferably not more than 0.3 탆, more preferably not more than 0.1 탆, in order to effectively prevent the TiN particles from becoming a fracture origin and further improve the impact resistance of the high strength steel sheet.

The average particle diameter of the TiN particles is obtained, for example, by the following method. That is, a sample for a transmission electron microscope (TEM) containing TiN particles was prepared from the plate thickness cross-section parallel to the rolling direction by an extraction replica method, and 10 or more TiN particles were observed using a transmission electron microscope. The particle diameter of each TiN particle is determined by the diameter of the circle having an area equivalent to the projected area of the TiN particle obtained by the image analysis. Then, the particle diameters of at least 10 TiN particles are measured, and the average particle diameter of the TiN particles is determined from the average value.

"AlN Particles"

In the steel sheet structure of the high strength steel sheet of the present invention, the density of the AlN grains having a grain diameter of 1 탆 or more is 1.0 / mm 2 or less. The coarse AlN grains having a particle size of 1 탆 or more become the fracture origin. In the steel sheet structure of the high-strength steel sheet of the present invention, the density of the AlN grains having a particle diameter of 1 mu m or more is 1.0 / mm2 or less, so that fracture of the AlN grains is prevented. The density of the AlN particles having a particle diameter of 1 占 퐉 or more is preferably 0.5 / mm2 or less and more preferably 0.1 / mm2 or less in order to more effectively prevent destruction of the AlN particles.

The measurement of the average particle diameter of these TiN particles and the measurement of the density of the AlN particles having a particle diameter of 1 占 퐉 or more may be carried out at any plate thickness position among the steel plates except for the outermost surface of the steel plate. For example, as in the case of residual austenite or ferrite described later, it is preferable to measure the steel sheet at a position 1/8 to 3/8 in thickness as a typical steel sheet.

In the present invention, AlN particles having a particle diameter of 1 mu m or more means AlN particles having a circle equivalent diameter d of 1 mu m or more. The circle equivalent diameter d is the diameter of a circle having an area equal to the projected area S of the particles obtained by the image analysis, and is obtained by the following expression. d =? (4S /?)

The density of the AlN particles in the present invention is determined, for example, by the following method.

That is, an area of 10.0 mm 2 or more of the plate thickness cross-section parallel to the rolling direction was observed by using a Field Emission Scanning Electron Microscope (FE-SEM), and the number of AlN grains of 1 탆 or more was counted, Density is calculated. The components of the AlN particles can be confirmed by using an energy dispersive X-ray spectrometer attached to the FE-SEM.

The steel sheet structure of the high strength steel sheet of the present invention contains 1 to 8% of residual austenite in a volume fraction of 1/8 to 3/8 of the thickness centered on 1/4 of the sheet thickness, The average aspect ratio of the retained austenite is 2.0 or less and the amount of dissolved Mn in the retained austenite is 1.1 times or more of the average amount of Mn.

Further, it is preferable that the entire steel sheet structure contains 1 to 8% of retained austenite in a volume fraction. However, the metal structure in the range of 1/8 thickness to 3/8 thickness centering on 1/4 of the plate thickness of the steel plate represents the structure of the whole steel plate. Therefore, if the steel sheet contains 1 to 8% of retained austenite in a volume fraction of 1/8 to 3/8 of the thickness of the steel sheet, the volume percentage of the entire steel sheet is 1 to 8% Of the retained austenite. For this reason, in the present invention, the range of the volume fraction of retained austenite in the range of 1/8 thickness to 3/8 thickness of the base material steel sheet is defined.

Further, in the steel sheet structure of the high strength steel sheet of the present invention, in addition to the retained austenite, in the range of 1/8 thickness to 3/8 thickness centered on 1/4 of the sheet thickness, 10 to 75% At least one of bainitic ferrite and bainite in a total amount of 10 to 50% and tempering martensite in an amount of 5 to 50% in total, wherein pearlite is limited to 5% or less in volume fraction, and fresh martensite It is preferable that the volume of the site is limited to 15% or less. When the high-strength steel sheet of the present invention has such a steel sheet structure, it has excellent moldability.

Likewise, it is preferable that these metal structures such as ferrite have a predetermined range in the entire steel sheet structure. However, the metal structure in the range of 1/8 thickness to 3/8 thickness centering on 1/4 of the plate thickness of the steel plate represents the structure of the whole steel plate. Therefore, in the range of 1/8 thickness to 3/8 thickness of the steel sheet, the volume fraction of the ferrite of 10 to 75% or less, the total of 10 to 50% of bainitic ferrite and bainite, If the pearlite is limited to 5% or less in the volume fraction and the fresh martensite is limited to 15% or less in the volume fraction, it is preferable that in the entire structure of the steel sheet, Ferrite and the like can be considered to be within a predetermined range. Therefore, in the present invention, the range of the volume fraction of the metal structure such as ferrite is specified within the range of 1/8 to 3/8 of the thickness of the steel sheet.

&Quot; Residual austenite "

The retained austenite needs to contain the retained austenite in a range that does not deteriorate the value of the elongation at low temperature in order to greatly improve the strength and ductility. When the volume percentage of retained austenite is less than 1%, the improvement in strength and ductility is insufficient, and this is made the lower limit. From the viewpoints of strength and moldability, the amount of retained austenite is preferably 1.5% or more, more preferably 2.0% or more. On the other hand, the retained austenite becomes a breaking point and greatly deteriorates the bendability. Therefore, it is necessary to limit the steel sheet structure to a volume fraction of 8% or less. In order to improve the bendability, it is more preferable that the volume fraction of the retained austenite is 6% or less.

Further, in order to prevent breakage of the retained austenite as a starting point, it is preferable that the retained austenite has a stable shape and is chemically stable.

In the present invention, the retained austenite has an average aspect ratio of 2.0 or less and has a stable shape with excellent isotropy. In order to make the shape of the retained austenite more stable, the average aspect ratio of the retained austenite is preferably 1.8 or less, more preferably 1.6 or less. The lower limit of the average aspect ratio of the retained austenite is 1.0. If the average aspect ratio exceeds 2.0, a part of the retained austenite easily transforms into martensite when stretched at a low temperature, and a fracture origin is generated, and the elongation aberration is deteriorated.

In the present invention, the amount of dissolved Mn in the retained austenite is 1.1 times or more the "average amount of Mn in the retained austenite / the average amount of Mn) ≥1.1", and the retained austenite is chemically stable. In order to chemically stabilize the retained austenite, the amount of Mn incorporated in the retained austenite is preferably 1.2 times or more of the average Mn amount, and more preferably 1.3 times or more. The upper limit is not specially set, but in order to make it 2.0 times or more, special equipment is required and the actual upper limit is 2.0 times.

"ferrite"

The ferrite is a structure effective for improving the shrinkage value at low temperature, and it is preferable that the ferrite is contained in the steel sheet structure in a volume fraction of 10 to 75%. If the volume fraction of ferrite is less than 10%, there is a possibility that a sufficient elongation value may not be obtained. The volume fraction of ferrite contained in the steel sheet structure is more preferably at least 15%, more preferably at least 20% from the viewpoint of the elongation value. On the other hand, since ferrite is a soft structure, when the volume fraction exceeds 75%, sufficient strength may not be obtained. In order to sufficiently increase the tensile strength of the steel sheet, the volume fraction of ferrite contained in the steel sheet structure is preferably 65% or less, more preferably 50% or less.

"Pearlite"

If the number of pearlite increases, the ductility deteriorates. From this point of view, the volume fraction of pearlite contained in the structure of the steel sheet is preferably limited to 5% or less, and more preferably 2% or less.

&Quot; Bainitic ferrite, bainite "

It is preferable that the bainitic ferrite and the bainite have a good balance of strength and ductility, and that the steel sheet structure contains either or both of bainitic ferrite and bainite having a total volume fraction of 10 to 50%. Bainitic ferrite and bainite are microstructures having medium strength between soft ferrite and hard martensite, tempered martensite and retained austenite, more preferably at least 15% from the viewpoint of stretch flangeability, More preferably 20% or more. On the other hand, if the volume fraction of bainitic ferrite and bainite exceeds 50% in total, the yield stress becomes excessively high and the shape crystallinity deteriorates, which is not preferable. The bainitic ferrite and bainite may contain either one or both of them.

"Fresh martensite"

Fresh martensite greatly improves the tensile strength, but on the other hand, it becomes a fracture origin and greatly deteriorates the value of the elongation at low temperature, so that it is preferable that the volume fraction is limited to 15% or less in the steel sheet structure. The volume fraction of fresh martensite is more preferably 10% or less, more preferably 5% or less, in order to increase the value of the elongation at low temperature.

"Tempering Martensite"

The tempering martensite is a structure that greatly improves the tensile strength, and may be contained in the steel sheet structure in a volume fraction of 50% or less. From the viewpoint of tensile strength, the volume fraction of the tempering martensite is preferably 10% or more. On the other hand, if the volume fraction of the tempered martensite contained in the steel sheet structure exceeds 50%, the yield stress becomes excessively high and the shape crystallinity deteriorates, which is not preferable.

"Other"

The steel sheet structure of the high-strength steel sheet of the present invention may contain a structure other than the above, such as coarse cementite. However, when the amount of coarse cementite is increased in the steel sheet structure, the bending property is deteriorated. From this point of view, the volume fraction of coarse cementite contained in the steel sheet structure is preferably 10% or less, more preferably 5% or less.

The volume fractions of the respective tissues contained in the steel sheet structure of the high-strength steel sheet of the present invention can be measured, for example, by the following methods.

The volume fraction of retained austenite was determined by performing an X-ray diffraction test on an arbitrary surface in a range of 1/8 thickness to 3/8 thickness parallel to the surface of the steel sheet to calculate the area fraction of the retained austenite, It can be regarded as a volume fraction in the range of 1/8 thickness to 3/8 thickness.

In addition, the microstructure in the 1/8 to 3/8 thickness range is highly homogeneous, and if a sufficiently large area is measured, 1/8 to 3/8 of the thickness, / 8 < / RTI > thickness is obtained. Concretely, it is preferable to carry out the X-ray diffraction test in a range of 250000 square m or more on a 1/4 thickness surface parallel to the plate surface of the steel sheet.

The fraction of the microstructure (ferrite, bainitic ferrite, bainite, tempered martensite, pearlite, fresh martensite) except for the retained austenite was observed with an electron microscope in the range of 1/8 thickness to 3/8 thickness . Concretely, a sample is taken as an observation surface perpendicular to the plate surface of the base material steel sheet and parallel to the rolling direction (downward direction), and the observation surface is polished or etched away. Then, the area fraction is measured by observing a range of 1/8 thickness to 3/8 thickness centering on 1/4 of the plate thickness with a Field Emission Scanning Electron Microscope (FE-SEM). In this case, for example, in the range of 1/8 thickness to 3/8 thickness, observation with an electron microscope is performed in three or more fields set at intervals of 1 mm or more. Then, the area fraction of each texture such as ferrite in the range of the observation area in the range of 5000 square m or more is calculated, and it is regarded as the volume fraction of each tissue in the range of 1/8 thickness to 3/8 thickness .

Ferrite is a crystal grain of a mass, and is an area in which there is no iron carbide having a long diameter of 100 nm or more inside. The volume fraction of ferrite is the sum of the volume fraction of the ferrite remaining at the maximum heating temperature and the volume fraction of the ferrite newly generated in the ferrite transformation temperature region.

The bainitic ferrite is a set of crystal grains in the form of lasers, and does not contain an iron-based carbide having a long diameter of 20 nm or more inside the lath.

Bainite is a set of crystal grains in the form of lasers and has a plurality of iron-based carbides having a long diameter of 20 nm or more in the inside of the lath, and further belongs to a group in which the carbides are single, i.e., an iron-based carbide group elongated in the same direction. Here, the iron-based carbide group extending in the same direction means that the difference in the elongation direction of the iron-based carbide group is within 5 degrees.

The tempering martensite is a set of crystal grains in the form of lasers and contains a plurality of iron-based carbides having a long diameter of 20 nm or more in the interior of the lath. Further, these carbides belong to a plurality of valleys, will be.

Further, by observing the iron-based carbide in the lase-shaped crystal grains by using the FE-SEM and examining the elongation direction thereof, the bainite and the tempering martensite can be easily distinguished.

Further, fresh martensite and retained austenite are not sufficiently corroded by the batt etching. Therefore, it is clearly distinguished from the above-described structure (ferrite, bainitic ferrite, bainite, tempered martensite) in observation by FE-SEM.

Therefore, the volume fraction of fresh martensite is obtained as the difference between the area fraction of the un-corroded area observed by the FE-SEM and the area fraction of the residual austenite measured by X-ray.

(Zinc plated layer)

Further, in the present invention, a high-strength galvanized steel sheet excellent in impact resistance characteristics comprising a zinc plating layer formed on the surface of a high-strength steel sheet can be obtained. The zinc plated layer may be alloyed. When the zinc plating layer is formed on the surface of the high-strength steel sheet, it has excellent corrosion resistance. In addition, when a galvanized zinc alloy layer is formed on the surface of the high-strength steel sheet, it has excellent corrosion resistance and excellent adhesion of the coating. The zinc plating layer or the galvanized zinc plating layer may contain Al as an impurity.

The alloyed zinc plated layer may contain one or more of Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca, Cu, Li, Ti, Be, Bi, Sr, I, Cs, Or they may be incorporated. The effect of the present invention is not impaired even if the alloying zinc plating layer contains one or more kinds of the above elements, or even if it is incorporated. In some cases, the corrosion resistance and workability are improved depending on the content.

The adhesion amount of the zinc plating layer or the zinc alloy plating layer is not particularly limited, but is preferably 20 g / m 2 or more from the viewpoint of corrosion resistance and 150 g / m 2 or less from the viewpoint of economy. The average thickness of the zinc-plated layer or the zinc-plated zinc-plated layer is not less than 1.0 탆 and not more than 50 탆. When the thickness is less than 1.0 탆, sufficient corrosion resistance is not obtained. Preferably 2.0 mu m or more. On the other hand, when it exceeds 50.0 탆, it is not economical and it damages the strength of the steel sheet. From the viewpoint of raw material cost, the thinner the thickness of the zinc plated layer or the galvanized layer is, the more preferable it is 30.0 占 퐉 or less.

The average thickness of the plated layer was measured by FE-SEM after finishing the plate thickness cross-section parallel to the rolling direction of the steel sheet and measuring the thickness of the plated layer in total of 10 places on each of the front and back surfaces of the steel sheet And the average value is taken as the thickness of the plating layer.

Further, in the alloying treatment, the content of iron in the galvannealed layer is preferably 8.0% or more and 9.0% or more in order to ensure good anti-flaking properties. The content of iron in the galvannealed layer is preferably 12.0% or less and 11.0% or less in order to ensure good resistance to powdering.

Further, in the present invention, a coating film containing a composite oxide containing phosphorus oxide and / or phosphorus may be formed on the surface of the zinc-plated layer or the alloyed zinc-plated layer. The coating film comprising the phosphorus oxide and / or the phosphorus-containing composite oxide can function as a lubricant when the steel sheet is processed, and the zinc plated layer formed on the surface of the steel sheet can be protected.

(Manufacturing method)

Next, a method for producing the high-strength steel sheet of the present invention will be described in detail.

In order to produce the high strength steel sheet of the present invention, a slab having the chemical composition (composition) is first cast.

As the slab to be provided for the hot rolling, a continuous cast slab or a thin slab castor can be used. The method of manufacturing a high strength steel sheet of the present invention is suitable for a process such as continuous casting-direct rolling (CC-DR) in which hot rolling is immediately performed after casting.

(Hot rolling process)

In the hot rolling step, the slab heating temperature needs to be 1210 占 폚 or higher in order to sufficiently dissolve the Ti-based inclusions generated at the time of casting and uniformly solidify Ti in the steel, and it is preferably 1225 占 폚 or higher . In addition, if the slab heating temperature is excessively low, the finish rolling temperature falls below the Ar ₃ transformation point. As a result, rolling is performed in the two-phase region of ferrite and austenite, and the hot rolled steel sheet structure becomes heterogeneous mixed grain structure. Even if the cold rolling process and the continuous annealing process are performed, the heterogeneous structure is not solved, Resulting in a steel sheet having poor bendability. Further, the lowering of the slab heating temperature leads to an increase in the excessive rolling load, which makes it difficult to carry out rolling, or may cause a defect in the shape of the steel sheet after rolling. The effect of the present invention is exhibited even if the upper limit of the slab heating temperature is not particularly specified. However, it is not preferable from the economical point of view to make the heating temperature excessively high, so the upper limit of the slab heating temperature is preferably 1350 캜 or lower.

The Ar ₃ transformation point is calculated by the following equation.

_{Ar 3 = 901-325 × C + 33} × Si-92 × (Mn + Ni / 2 + Cr / 2 + Cu / 2 + Mo / 2) + 52 × Al

In the above formula, C, Si, Mn, Ni, Cr, Cu, Mo, and Al are the content [mass%] of each element. The element not contained is calculated as 0.

In the present invention, after the slab is heated to the above-mentioned heating temperature, the temperature is lowered in the range of at least 1100 to 1000 ° C under a condition satisfying the following formula (1). Ti represents the machining temperature of the i-th pass, ti represents the elapsed time from the i-th pass to the (i + 1) th pass, and ei represents the reduction ratio of the i-th pass.

In order to inhibit the formation of coarse Ti nitride or Al nitride and to produce a steel sheet containing fine TiN particles, it is necessary to heat a large amount of the potential of the Ti nitride formation site in the steel by hot rolling in the temperature range of 1100 to 1000 ° C . However, at a temperature range of 1100 to 1000 占 폚, the dislocations introduced by machining easily disappear due to diffusion of Fe atoms. As a result, it is necessary to continuously carry out the machining (pressing down) in which a deformation amount is obtained so that the potential is sufficiently introduced, in a relatively short time. In other words, it is necessary to control the number of passes to a plurality of times, to shorten the elapsed time between adjacent passes, and to appropriately control the processing temperature and the reduction rate in each of the passes.

In the hot rolling process, it is in the temperature range of after taking out the slab from heating to, or up to 850 ℃ rolling completion temperature to the Ar ₃ temperature side is high as the lower limit may be the reduction of any of the number of passes. Since the hot rolled steel sheet strongly affects the dispersed state of the TiN and AlN grains in the hot rolled steel sheet in the range of 1100 DEG C to 1000 DEG C, the hot rolling conditions in the same temperature range And

The potential introduced at the time of pressing and deforming performed in the temperature range exceeding 1100 占 폚 is immediately extinguished and does not act as a precipitation site of TiN and does not affect the dispersion state of TiN and AlN particles which are problematic. On the other hand, nucleation of the particles which can be coarse TiN and AlN is completed until rolling is carried out in the range of less than 1000 DEG C, and rolling after that (temperature range of less than 1000 DEG C) It does not affect the dispersion state of the particles.

Generally, rolling is performed in 8 to 25 passes between the heating furnace and the completion of rolling. And is performed in the range of 1100 DEG C to 1000 DEG C in 2 to 10 passes. The thickness of the sheet to be pressed in the temperature range starts from 200 to 500 mm, and is generally rolled to 10 to 50 mm. The plate width is generally 500 to 2000 mm. The temperature of the steel sheet is the temperature of the surface of the steel sheet, and the measurement method thereof is not particularly limited. However, it may be measured directly using, for example, a thermocouple.

Specifically, in Expression (1), for example, the number of passes i may be in the range of 2 to 10, preferably in the range of 5 to 8. [ The elapsed time from the i-th path to the (i + 1) th path may be in the range of 2 to 300 seconds, preferably in the range of 5 to 180 seconds, and more preferably in the range of 10 to 120 seconds.

The processing temperature of the first pass, which is the first pass in the hot rolling in the temperature range of 1100 to 1000 占 폚, may be in the range of 1100 to 1050 占 폚, preferably 1090 to 1065 占 폚. The reduction ratio in the i-th pass may be in the range of 5 to 50%, preferably in the range of 15 to 35%.

(Equation 1) is an empirical equation showing the generation behavior of TiN particles, and is a product of the polynomial expression representing the driving force of particle generation, the exponent of the atomic diffusion coefficient and the time t, The amount of dislocations is represented by the strain amount?, And is multiplied by them. If the value shown in the formula (1) is less than 1.0, the production of TiN becomes insufficient, and the solid solution N remains until the hot rolling to 1000 deg. C, and coarse AlN is produced. On the other hand, if the value shown in the formula (1) exceeds 5.0, the generation of TiN becomes excessively active, and the coarsening of TiN progresses, thereby deteriorating the barrier property.

In the present invention, the elapsed time between a plurality of adjacent paths is controlled in a relatively short time by performing the depressing under a condition satisfying the above-mentioned (formula 1) in a temperature range of at least 1100 to 1000 ° C, Since the processing temperature and the reduction rate in the pass are appropriately controlled, it is possible to introduce a large amount of the potential, which is the site of production of the Ti nitride, into the steel, and to produce fine Ti nitride in the steel. Further, there is no particular limitation on the reduction carried out in the temperature range exceeding 1100 占 폚 and the reduction carried out in the temperature range lower than 1000 占 폚. For example, the pressing may be performed under a condition satisfying the above-mentioned (formula 1) in a temperature range exceeding 1100 占 폚, or may be carried out under a condition not satisfying the above formula (1). Alternatively, it is not necessary to perform the pressing down in a temperature range exceeding 1100 占 폚. Similarly, the pressing may be carried out under a condition satisfying the above-mentioned (formula 1) in a temperature range of less than 1000 ° C, or may be carried out under a condition not satisfying the above formula (1).

In the present invention, at least 1100 to a temperature range of 1000 ℃, after carrying out the hot rolling under the condition satisfying the above equation (1), 800 ℃ and the Ar ₃ transformation point or the finish hot rolling of more than 970 ℃ than the temperature of the high-side of the Completion at temperature and winding in the temperature range of 750 ° C or less. The plate thickness after finish rolling is, for example, 2 mm to 10 mm. If the finishing rolling temperature is less than 800 占 폚, the rolling load at the finish rolling becomes high, which makes it difficult to perform the hot rolling or may lead to a defective shape of the hot rolled steel sheet obtained after hot rolling. If the finishing rolling temperature is lower than the Ar ₃ transformation point, the hot rolling may cause rolling of the two-phase region of ferrite and austenite, and the hot-rolled steel sheet may have an uneven homogeneous structure. On the other hand, if the upper limit of the finishing rolling temperature is 970 캜 or higher, the production of TiN becomes insufficient, and the remaining N may possibly generate Al and nitride.

In the present invention, in the hot rolling process, of less than 1100 to over a temperature range of 1000 ℃ subjected to hot rolling under the condition satisfying the above equation (1), at least the higher of 800 ℃ and the Ar ₃ transformation point temperature 970 ℃ It is possible to inhibit the formation of coarse Ti nitrides in the temperature range of 1100 to 1000 占 폚 and to form fine TiN grains between 1000 占 폚 and the finish hot rolling temperature. As a result, the finally obtained high strength steel sheet has excellent impact resistance characteristics.

In order to prevent excessive increase in the thickness of the oxide formed on the surface of the hot-rolled steel sheet and deterioration of the pickling ability, the coiling temperature is set to 750 캜 or lower. In order to further increase the pickling ability, the coiling temperature is preferably 720 占 폚 or lower, more preferably 700 占 폚 or lower.

On the other hand, if the coiling temperature is less than 500 캜, the strength of the hot-rolled steel sheet becomes excessively high, and cold rolling becomes difficult, so that the coiling temperature is preferably 500 캜 or higher. In order to reduce the load of cold rolling, the coiling temperature is preferably 550 DEG C or higher, more preferably 600 DEG C or higher.

Then, the hot rolled steel sheet wound in the temperature region is cooled at an average cooling rate of 15 DEG C / hour or less. As a result, the distribution of Mn dissolved in the steel sheet proceeds and the retained austenite can selectively remain in the concentrated region of Mn, and the amount of Mn incorporated in the retained austenite can be increased. As a result, in the finally obtained high strength steel sheet, the amount of dissolved Mn in the retained austenite is 1.1 times or more of the average amount of Mn. The distribution of Mn after winding is likely to proceed at higher temperatures. Therefore, in particular, in the range from the coiling temperature (coiling temperature -50 占 폚), the cooling rate of the steel sheet needs to be 15 占 폚 / hour or less.

Then, the hot-rolled steel sheet thus produced is preferably subjected to pickling. Since pickling removes the surface oxide of the hot-rolled steel sheet, it is important to improve the plating performance of the steel sheet. The pickling may be performed once or divided into a plurality of steps.

(Cold rolling process)

Subsequently, the hot rolled steel sheet after pickling is subjected to a cold rolling step of cold rolling at a reduction ratio of 30 to 75% in order to make the retained austenite have a stable shape excellent in isotropy. If the reduction ratio is less than 30%, the shape of the retained austenite can not be stabilized, and the final high strength steel sheet can not have an average aspect ratio of retained austenite of 2.0 or less. In order to make the retained austenite have a stable shape, the reduction ratio in the cold rolling step is preferably 40% or more, more preferably 45% or more. On the other hand, in cold rolling in which the reduction rate exceeds 75%, the cold rolling load becomes too large and cold rolling becomes difficult. From this point of view, the reduction rate is preferably 75% or less. From the viewpoint of cold rolling load, it is more preferable that the reduction rate is 70% or less.

In the cold rolling step, the effects of the present invention are exhibited even when the number of rolling passes and the reduction rate for each rolling pass are not particularly specified.

(Continuous annealing process)

Then, the cold-rolled steel sheet obtained after the cold-rolling step is passed through a continuous annealing line to carry out a continuous annealing process. In the continuous annealing step of the present invention, the temperature range of 550 to 700 占 폚 is heated at an average heating rate of 10 占 폚 / sec or less, the maximum heating temperature is set between (Ac ₁ transformation point +40) Cooling at an average cooling rate of 1.0 to 10.0 ° C / sec in a temperature range of from 700 to 500 ° C, cooling at an average cooling rate of 5.0 to 200.0 ° C / sec in a temperature range of 700 to 500 ° C, For 30 to 1000 seconds. Thus, the high strength steel sheet of the present invention is obtained.

In the continuous annealing step, the temperature range of 550 to 700 占 폚 is heated at an average heating rate of 10 占 폚 / sec or less so that the recrystallization of the cold-rolled steel sheet proceeds sufficiently and the shape of the retained austenite becomes more stable, Finally, the remaining austenite has a shape close to the sphere shape. If the average heating rate in the temperature range of 550 to 700 占 폚 exceeds 10 占 폚 / sec, the shape of the retained austenite can not be stabilized.

Further, in the maximum heating temperature (Ac ₁ transformation point + 40) under ℃ in the continuous annealing process, since the iron-based carbide coarse the steel sheet number remaining melt, formability is remarkably deteriorated, and the maximum heating temperature (Ac ₁ transformation point + 40) ℃ or more. From the viewpoint of moldability, the maximum heating temperature is preferably not less than (Ac ₁ transformation point + 50) ° C, and more preferably not less than (Ac ₁ transformation point + 60) ° C. On the other hand, when the maximum heating temperature exceeds 1000 ° C, the diffusion of atoms is promoted and the distribution of Si, Mn and Al becomes weak. Therefore, the maximum heating temperature is set to 1000 ° C or lower. In order to control the amounts of Si, Mn and Al in the retained austenite, the maximum heating temperature is preferably not higher than the Ac ₃ transformation point temperature.

If the average cooling rate exceeds 10.0 DEG C / second in the temperature range from the maximum heating temperature to 700 DEG C, the ferrite fraction in the steel sheet tends to be uneven and the formability is deteriorated, so that the upper limit of the average cooling rate is set at 10.0 DEG C / second . On the other hand, when the average cooling rate is less than 1.0 占 폚 / sec, a large amount of ferrite and pearlite are produced and the retained austenite is not obtained, so the lower limit of the average cooling rate is 1.0 占 폚 / sec. In order to obtain the retained austenite, the average cooling rate is preferably 2.0 DEG C / second or more, more preferably 3.0 DEG C / second or more.

When the average cooling rate is less than 5.0 占 폚 / sec in the temperature range of 700 to 500 占 폚, a large amount of pearlite and / or iron-based carbide is generated and no residual austenite remains. Or more. From this viewpoint, the average cooling rate is preferably 7.0 DEG C / second or more, and more preferably 8.0 DEG C / second or more. On the other hand, the effect of the present invention is exerted even if the upper limit of the average cooling rate is not particularly specified. However, in order to exceed the average cooling rate of 200 DEG C / second, special facilities are required, Second.

Further, in order to proceed the bainite transformation and obtain the retained austenite, a rectification treatment is performed in a temperature range of 350 to 450 DEG C for 30 to 1000 seconds. If the rectification time is short, the bainite transformation does not proceed, the concentration of C in the retained austenite becomes insufficient, and the retained austenite can not be sufficiently left. From this point of view, the lower limit of the rectification time is 30 seconds. The rectification time is preferably 40 seconds or more, more preferably 60 seconds or more. On the other hand, if the rectification time is excessively long, the iron-based carbide is produced, C is consumed as the iron-based carbide, and the retained austenite is not sufficiently obtained. From this viewpoint, the rectification time is preferably 800 seconds or less, and more preferably 600 seconds or less.

Further, in the present invention, in the continuous annealing step of the above-described manufacturing method, galvanization may be performed after the rectifying treatment to form a zinc-plated layer on the surface of the steel sheet to form a high-strength galvanized steel sheet.

Further, in the present invention, in the continuous annealing step of the above-described production method, after cooling in the temperature range of 700 to 500 占 폚, or before or after the rectifying treatment in the temperature range of 350 to 450 占 폚, The steel sheet may be immersed in a plating bath to form a galvanized layer on the surface of the steel sheet to form a high strength galvanized steel sheet.

Thereby, a high-strength galvanized steel sheet excellent in impact resistance properties having a zinc plated layer formed on its surface can be obtained.

The zinc plating bath is not particularly limited and may be a zinc plating bath containing at least one of Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca, Cu, Li, Ti, , REM, or the like are mixed, the effect of the present invention is not impaired, and depending on the amount, the corrosion resistance and workability are improved. Further, Al may be contained in the zinc plating bath. In this case, the Al concentration in the bath is preferably 0.05% or more and 0.15% or less.

The temperature of the alloying treatment is preferably 480 to 560 占 폚, and the residence time of the alloying treatment is preferably 15 to 60 seconds.

Alternatively, after the steel sheet is immersed in a zinc plating bath, the steel sheet may be reheated to 460 DEG C to 600 DEG C and held for 2 seconds or longer to alloy the zinc plating layer.

By performing such an alloying treatment, a Zn-Fe alloy in which a zinc plating layer is alloyed is formed on the surface, and a high strength galvanized steel sheet having a galvanized layer alloyed on its surface is obtained.

It is also possible to provide a coating film containing a composite oxide containing phosphorus oxide and / or phosphorus on the surface of the galvanized layer or alloyed zinc plating layer of the high-strength galvanized steel sheet.

In the present embodiment, it is preferable that the alloying treatment is carried out at a temperature of 200 to 350 DEG C for 30 to 1000 seconds. Thereby, the steel sheet structure contains tempering martensite.

After the alloying treatment, the steel sheet after the alloying treatment is cooled to 350 占 폚 or less to produce martensite, and then heated at a temperature in the range of 350 占 폚 to 550 占 폚 , And is allowed to stand for at least 2 seconds to produce tempered martensite. In the continuous annealing step, the steel sheet cooled to a temperature region of 500 DEG C or lower is further cooled to 350 DEG C or lower to generate martensite, and then reheated and retained at 400 SIMILAR 500 DEG C to obtain tempered martensite Is generated.

The present invention is not limited to the above examples.

For example, in order to improve the plating adhesion, the steel sheet before annealing may be plated with one or more selected from Ni, Cu, Co, and Fe.

Further, in the present embodiment, the steel sheet subjected to annealing may be temper rolling for the purpose of shape correction. However, when the reduction rate after annealing exceeds 10%, the soft ferrite portion undergoes work hardening and the ductility is greatly deteriorated, so that the reduction rate is preferably less than 10%.

<Examples>

The present invention will be described in more detail with reference to Examples.

(Composition) of A to AF shown in Tables 1 and 2 and the chemical components (compositions) of BA to BC shown in Table 3 were cast, and after the casting, that is, under the conditions shown in Tables 4 to 7 Rolled at a heating temperature, a rolling start temperature, a value of (Equation 1) in hot rolling at a temperature range of 1100 to 1000 캜, a hot rolling hot finish temperature, Cooled to an average cooling rate shown in Tables 4 to 7, and acid washed. Thereafter, cold rolling at the reduction ratio shown in Tables 4 to 7 was carried out.

Subsequently, annealing was carried out under the conditions shown in Tables 8 to 11 to obtain steel plates of Experimental Examples 1 to 108 and 201 to 208. In the annealing step, the temperature range of 550 to 700 占 폚 is heated at the average heating rate shown in Tables 6 to 8, heated to the maximum heating temperature shown in Tables 8 to 11, (Cooling rate 1) shown in Tables 8 to 11 and the temperature range of 700 to 500 占 폚 was cooled to the average cooling rate (cooling rate 2) shown in Tables 8 to 11, A rectifying process of rectifying the temperature range of 450 占 폚 to the times shown in Tables 8 to 11 was performed and then cooled to room temperature.

After cooling to room temperature, cold rolling of 0.15% was carried out in Experimental Examples 6 to 20 and Experimental Examples 70 to 108, cold rolling of 1.50% in Experimental Example 23, and cold rolling of 1.00% in Experimental Example 28 And in Examples 31 to 55, cold rolling of 0.25% was carried out.

In Examples 34, 44, 78, and 81, electroplating was performed by an electroplating line after the annealing process, thereby forming an electro galvanized steel sheet.

In Examples 19, 24 and 84, the steel sheet was cooled from a cooling rate of 2 to 500 ° C and then dipped in a zinc plating bath until the steel sheet was cooled to a temperature range of 350 to 450 ° C to obtain a hot-dip galvanized steel sheet.

In Examples 29 and 87, the steel sheet was immersed in a zinc plating bath after a rectifying treatment in a temperature range of 350 to 450 ° C, and then cooled to room temperature to obtain a hot-dip galvanized steel sheet.

In Experimental Examples 4, 14 and 75, the steel sheet was cooled from a cooling rate of 2 to 500 ° C and then immersed in a zinc plating bath until the steel sheet was cooled to a temperature range of 350 to 450 ° C. Further, The steel sheet was subjected to alloying treatment at a temperature of 30 seconds to obtain a galvannealed steel sheet.

In Experimental Examples 9, 58 and 72, alloying treatment was carried out by immersing in a zinc plating bath after the rectifying treatment in the temperature range of 350 to 450 ° C and further holding for 30 seconds at the alloying temperature shown in Tables 8 to 11 , And an alloyed hot-dip galvanized steel sheet.

In Examples 14 and 72, a coating film containing a composite oxide containing phosphorus was provided on the surface of the zinc plated layer.

&Quot; GI &quot; means a hot-dip galvanized steel sheet, &quot; EG &quot; denotes an electrodeposited zinc-coated galvanized steel sheet, Plated steel plate means.

The microstructure in the range of 1/8 thickness to 3/8 thickness in the steel sheets of Experimental Examples 1 to 108 and 201 to 208 was observed and the volume fraction was measured. The results are shown in Tables 12 to 15. In Table 12 to Table 15, "F" means ferrite, "B" means bainite, "BF" means bainitic ferrite, "TM" means tempering martensite, "M Quot; means fresh martensite, and &quot; residual y &quot; means residual austenite.

The amount of retained austenite in the microstructural fraction was measured by cutting out a plate thickness section and measuring the electron beam backscattering (EBSD) in a field-emission scanning electron microscope (FE-SEM) Electron Back Scattering Diffraction) analyzer, and the other section was polished to a specular surface and then etched away, and observed by FE-SEM.

The average aspect ratio (? Aspect ratio) of the retained austenite (?) Was obtained by measuring the aspect ratio of 20 retained austenites from the larger one in the map of the retained austenite obtained by the above-mentioned EBSD analyzer, And the EBSD analysis was carried out in the same manner to measure the aspect ratio of the 20 retained austenites from the larger side and the average value of the aspect ratios of the 40 retained austenites was calculated Respectively.

A sample for transmission electron microscope (TEM) was prepared by the extraction replica method from the aspect of observing the volume fraction of microstructure as the average particle diameter (TiN average size) of TiN particles, Circle-equivalent diameter) was measured, and the average value thereof was determined.

The inclusions in the range of 10.0 mm 2 were observed by FE-SEM on the side where the volume fraction of the microstructure was observed as the density of the AlN particles having a particle diameter of 1 탆 or more and the composition of inclusions having a circle equivalent diameter exceeding 1.0 탆 was measured , And the number of inclusions identified as AlN was counted to determine the density.

The ratio (WMn? / WMn) of the amount of Mn Mn in the retained austenite to the average Mn amount (WMn?) Was determined by measuring WMn and WMn? By the following method.

That is, EPMA analysis was performed in the same range as the EBSD analysis on the observation surface where the microstructural fraction was obtained, WMn was obtained from the obtained Mn concentration map, and furthermore, the Mn concentration map and the residual austenite map were superimposed. Only the measurement value of the Mn concentration was extracted, and WMn? Was obtained as the average value.

Tables 16 to 19 show the results of evaluation of the properties of the steel sheets of Experimental Examples 1 to 108 and 201 to 208 by the following methods.

Tensile test specimens in accordance with JIS Z 2201 were taken from the steel sheets of Experimental Examples 1 to 108 and 201 to 208 in accordance with JIS Z 2241 and tensile test was carried out in accordance with JIS Z 2241 to determine the yield stress "YS", the tensile strength "TS" EL &quot;

Further, a hole expansion test (JFST1001) for evaluating the flangeability was carried out to calculate the hole expansion limit value &quot;? &Quot;, which is an index of extension flangeability.

Further, the same tensile test piece was immersed in an alcohol to which liquid nitrogen had been added, cooled to -60 캜, taken out, and immediately subjected to a tensile test to determine the elongation at break (elongation value).

As shown in Tables 16 to 19, all the Experimental Examples 1 to 108 and 201 to 208 of Examples of the present invention all had a tensile strength of 900 MPa or more and a high shrinkage value of 20% or more, The characteristics were excellent.

On the other hand, the experimental examples of the comparative examples of Experimental Examples 1 to 108 were not high in tensile strength and excellent in impact resistance because the tensile strength was less than 900 MPa and / or the result of the elongation value was low.

Further, in Experimental Examples 14 and 72, a coating film containing a composite oxide containing phosphorus was provided on the surface of the zinc plated layer, and good characteristics were obtained.

In Experimental Example 5, the heating temperature of the slab before the hot rolling was low, coarse TiN remained, and the elongation at low temperature was low.

In Experimental Example 10, there is a large value of (Equation 1), a coarse TiN exists, and in Experimental Example 59, the value of (Equation 1) is small and coarse AlN exists. In Experimental Example 10 and Experimental Example 59, the decreasing value at low temperature is a dullness.

Experimental Example 15 is an example in which the hot rolling temperature of hot rolling is low and the microstructure is inhomogeneous in one direction, so that ductility values at ductility, elongation flangeability and low temperature are inferior.

Experimental Example 20 shows a high degree of coiling after hot rolling, and microstructures are very coarse, so that ductility values at ductility, stretch flangeability and low temperature are inferior.

Experimental Example 25 shows that the average cooling rate after winding is high, WMn? / WMn is low, Mn enrichment to retained austenite is insufficient, and the elongation at low temperature is inferior.

In Experimental Example 30, since the reduction rate of cold rolling is small and the aspect ratio (? Aspect ratio) of the retained austenite is large, the reduction value at low temperature is inferior.

In Experimental Example 35, the average heating rate of the annealing is large and the aspect ratio of the retained austenite (gamma aspect ratio) is large.

Experimental Example 40 is an example in which the maximum heating temperature in annealing is low and contains a large number of coarse iron carbide which is a starting point of fracture, so that ductility values at ductility, elongation flangeability and low temperature are inferior.

In Experimental Example 45, the cooling rate to 700 ° C was excessively high, and sufficient soft tissues were not obtained, so that the elongation at ductility and at low temperature was inferior.

In Experimental Example 50, the cooling rate 1 is excessively low, coarse carbides are produced, hard tissues are not sufficiently obtained, and the strength values are inferior in elongation, ductility, stretch flangeability and low temperature.

In Experimental Example 54, the rectification time at 350 to 450 ° C was short, the retained austenite was small, and the elongation at the ductile and low temperature was low.

In Experimental Example 55, the retention time at 350 to 450 캜 was long, the retained austenite was small and coarse carbide was produced, and the elongation at the ductility and the low temperature was inferior.

In Experimental Example 60, cooling rate 2 is low and coarse carbide is produced, so that ductility values at ductility, elongation flangeability and low temperature are inferior.

Experimental Examples 103 to 108 are examples in which the chemical components deviate from a predetermined range, and all of them can not obtain an elongation value at a sufficiently low temperature.

Claims (14)

In terms of% by mass,
C: 0.075 to 0.300%
Si: 0.30 to 2.50%
Mn: 1.30 to 3.50%
P: 0.001 to 0.050%,
S: 0.0001 to 0.0050%,
Al: 0.001 to 0.050%
Ti: 0.0010 to 0.0150%,
N: 0.0001 to 0.0050%,
O: 0.0001 to 0.0030%
, The balance being iron and inevitable impurities,
In the range of 1/8 thickness to 3/8 thickness centered on 1/4 of the plate thickness, contains 1 to 8% of retained austenite in a volume fraction and the average aspect ratio of the retained austenite is 2.0 or less , And the amount of dissolved Mn in the retained austenite is 1.1 times or more of the average amount of Mn,
A steel sheet structure containing TiN particles having an average particle diameter of 0.5 占 퐉 or less and having a density of AlN particles having a particle diameter of 1 占 퐉 or more of 1.0 pieces /
And a maximum tensile strength of 900 MPa or more. The steel sheet according to claim 1, wherein the steel sheet structure has a thickness of 1/8 to 3/8 the thickness of the base steel sheet, wherein the steel sheet has a volume fraction of 10 to 75% of ferrite and 10 to 50% of bainitic ferrite And either or both of bainite and 10 to 50% of tempering martensite,
A high strength steel sheet excellent in impact resistance property in which pearlite is limited to 5% or less in volume fraction and fresh martensite is limited to 15% or less in volume fraction. The method according to claim 1, further comprising at least one of Nb, V, B, Cr, Ni, Cu, Mo, W, Ca, Mg, Zr, Hf, REM,
The content of each element is expressed in mass%
Nb: 0.0010 to 0.0150%,
V: 0.010 to 0.150%,
B: 0.0001 to 0.0100%,
Cr: 0.01 to 2.00%
Ni: 0.01 to 2.00%
Cu: 0.01 to 2.00%
Mo: 0.01 to 1.00%
W: 0.01 to 1.00%
Ca, Mg, Zr, Hf and REM: 0.0001 to 0.5000% in total. delete delete The high strength galvanized steel sheet according to claim 1, wherein a galvanized layer is formed on the surface thereof, and is excellent in impact resistance. The high strength galvanized steel sheet according to claim 6, wherein a coating film is formed on the surface of the zinc plated layer, the coating comprising a composite oxide containing phosphorus oxide and / or phosphorus. In terms of% by mass,
C: 0.075 to 0.300%
Si: 0.30 to 2.50%
Mn: 1.30 to 3.50%
P: 0.001 to 0.050%,
S: 0.0001 to 0.0050%,
Al: 0.001 to 0.050%
Ti: 0.0010 to 0.0150%,
N: 0.0001 to 0.0050%,
O: 0.0001 to 0.0030%
And the remaining amount of iron and inevitable impurities is heated to 1210 占 폚 or more and the temperature is in the range of at least 1100 to 1000 占 폚, ℃ and hot rolling step of Ar ₃ or more temperature the higher of the transformation point to complete the rolling in the finishing hot rolling temperature below 970 ℃, and wound in a temperature range of less than 750 ℃, cooling at an average cooling rate of less when 15 ℃ / and ,
A cold rolling step of cold rolling at a reduction ratio of 30 to 75% after the hot rolling step,
After the cold rolling step, the temperature range of 550 to 700 占 폚 is heated at an average heating rate of 10 占 폚 / sec or less, the maximum heating temperature is between (Ac ₁ transformation point +40) Cooling at an average cooling rate of 1.0 to 10.0 占 폚 / sec in a temperature range, cooling at an average cooling rate of 5.0 to 200.0 占 폚 / sec in a temperature range of 700 to 500 占 폚, and cooling at a temperature range of 350 to 450 占 폚 of 30 to 1000 And a continuous annealing step of performing annealing for super-rectification treatment.
(1)

Ti represents the machining temperature of the i-th pass, ti represents the elapsed time from the i-th pass to the (i + 1) th pass, and ei represents the reduction ratio of the i-th pass. A process for producing a high strength galvanized steel sheet excellent in impact resistance, which comprises performing galvanization after rectifying treatment in a continuous annealing process according to claim 8 to form a zinc plating layer on the surface of the steel sheet. The steel sheet according to claim 8, wherein, in the continuous annealing step of the manufacturing method, after the steel sheet is cooled in the temperature range of 700 to 500 ° C, or before the rectifying treatment in the temperature range of 350 to 450 ° C, To form a galvanized layer on the surface of the steel sheet, wherein the galvanized steel sheet is excellent in impact resistance. The method of manufacturing a high strength galvanized steel sheet according to claim 10, wherein the galvanized steel sheet is immersed in the galvanizing bath, reheating the steel sheet to 460 to 600 캜, and holding the galvanized steel sheet for at least 2 seconds to alloy the galvanized layer. The method of manufacturing a zinc-plated steel sheet according to claim 10, further comprising, after forming the zinc-plated layer, coating a surface of the zinc-plated layer with a coating comprising a composite oxide containing phosphorus oxide and / or phosphorus A method of manufacturing a plated steel sheet. 12. The method of manufacturing a zinc-plated steel sheet according to claim 11, wherein the galvanizing layer is alloyed and then a coating film comprising a composite oxide containing phosphorus oxide and / or phosphorus is provided on the surface of the alloyed zinc- A method for manufacturing a high strength galvanized steel sheet.

The slab according to claim 8, wherein the slab further contains one or more of Nb, V, B, Cr, Ni, Cu, Mo, W, Ca, Mg, Zr, Hf,
The content of each element is expressed in mass%
Nb: 0.0010 to 0.0150%,
V: 0.010 to 0.150%,
B: 0.0001 to 0.0100%,
Cr: 0.01 to 2.00%
Ni: 0.01 to 2.00%
Cu: 0.01 to 2.00%
Mo: 0.01 to 1.00%
W: 0.01 to 1.00%
Ca, Mg, Zr, Hf and REM: 0.0001 to 0.5000% in total.